1. INTRODUCTION
FCC alloys have been extensively used in various industries, including nuclear plants,
aerospace, information technology, the automobile sector, and food industry, due to
their lightweight, favorable strength-to-weight ratio, good recyclability, low price,
and high corrosion resistance in different environments [1]. Aluminum (Al) and Copper (Cu) belong to the category of FCC materials and have
attracted significant attention across diverse industries, such as the electronics,
utensils, energy transportation, and electromechanical fields. Al alloys are typically
manufactured using either direct-chill (DC) casting or twin-roll casting (TRC)/continuous
casting. DC casting involves several steps, including casting, scalping, homogenization,
hot rolling (HR), room-temperature rolling (RTR), and annealing. In contrast, TRC
entails continuous casting, RTR, and annealing, which notably reduces production costs
[2].
This review discusses differences in the microstructure and texture evolution in DC
and TRC cast Al alloys, in subsequent sections. Due to the significant differences
in stacking fault energy (SFE) between Al and Cu alloys, the evolution of deformation-induced
products, such as deformation twins (DTs), strain localizations, and shear bands (SBs),
varies considerably in these two FCC materials. Al alloys have high SFE, while Cu
alloys are considered medium-low SFE materials. Pure Cu has an SFE of 78 mJ/m2, while the addition of Al, in amounts up to 2% and 4.5%, decreases the SFE of Cu
alloys to minimum values of 37 and 7 mJ/m2, respectively [3]. Materials with low SFE have large stacking fault (SF) regions that hinder dislocation
motion via slip during plastic deformation. However, the twinning mechanism is very
active in low SFE materials and facilitates further deformation. Plastic deformation
by twinning is reported to enhance the work-hardening rate while simultaneously increasing
strength and ductility [3]. Since Cu alloys are predominantly used in applications requiring high electrical
conductivity, achieving a trade-off between high conductivity and high strength is
vital. Thermo-mechanical processing (TMP) of Cu-Fe-P, Cu-Cr, Cu-Mg, Cu-Sn, and Cu-Fe
alloys strengthens the alloys through solid solution strengthening, work-hardening,
precipitation hardening, and also enhances the conductivity of the final Cu products
[4].
Further, this review discusses the evolution of microstructure in austenitic stainless
steels (ASSs) and high entropy alloys (HEAs). These alloys fall under the category
of FCC materials, with low SFE values of 18 mJ/m2 and 3.5 mJ/m2, respectively [5]. ASSs are principally divided into two series: the 300 series and the 200 series.
In the 200 series, nickel is replaced by nitrogen and/or manganese to stabilize the
austenitic structure [6]. Due to the rising cost of nickel, industries have pivoted towards low-nickel or
nickelfree ASSs (200 series). ASSs are widely employed in the fabrication of high-performance
pressure vessels and in medical/surgical tools due to their superior corrosion properties
and high formability [7]. However, their low yield strength limits their applications; this limitation can
be mitigated through TMP, strain hardening, and solid solution strengthening [8]. Additionally, since rolling of ASSs is an inevitable stage during the fabrication
process, the evolution of strain-induced martensite (α'), deformation bands, microbands,
and deformation twins (DTs) during rolling has been reported.
HEAs are alloys composed of multiple principal elements, with each element comprising
between 5% and 35% of the total composition [9]. Initially, their high configurational entropy was believed to be the sole factor
affecting phase stability. However, the role of mixing enthalpy was later found to
be equally critical for phase stabilization [10]. The equiatomic composition Co20Cr20Fe20Mn20Ni20, popularly known as the Cantor alloy, exists in a single FCC phase [11]. The SFE, estimated by density functional theory (DFT) calculations for the Cantor
alloy, is found to be approximately 20-25 mJ/m2, placing it in the low SFE regime [12]. Consequently, HEAs, which are classed as low SFE materials, characteristically
undergo twinning during deformation (a twinning-induced plasticity effect [13]) and subsequent annealing heat treatment. Deformation twinning (DT) incrementally
introduces more twin interfaces, which act as barriers (akin to grain boundaries (GBs))
to dislocation motion, and reduce the mean free path of dislocations, culminating
in the dynamic Hall-Petch effect [14]. The twinning-induced plasticity effect in HEAs, combined with the dynamic Hall-Petch
effect, results in higher ductility and strain hardening rates, giving them excellent
mechanical properties [15–17].
FCC metals/alloys are widely used in almost every sector, as highlighted above and
depicted in Fig 1. The requirement for good formability along with high strength, especially in the
creation of automotive frames, has been met by employing Al alloys and ASSs (Fig 1(a)). With the development of electric vehicles, the usage of lightweight materials such
as Al and Cu alloys has increased (Fig 1(b)). ASSs, known for their high corrosion resistance and strength, are used to manufacture
containers for the chemical, food, and paper industries (Fig 1(c)). HEAs have been developed to maintain mechanical strength at the high temperatures
found in airplane engines, thereby enhancing engine efficiency (Fig 1(d)). The fabrication of these alloys primarily consists of casting, rolling at high
and room temperatures, followed by intermediate annealing or solution annealing, etc.,
to create the final product for various applications. Predominantly, these processes
are used to produce metal sheets, and during the fabrication process, these alloys
often undergo the evolution of various microstructures and crystallographic textures.
The microstructure and texture of all metals/alloys have a strong relationship with
the material properties. For example, the balance between the deformation texture
components (Copper ({112}<111>), S {123}<634> and Brass ({110}<112>)) and annealing
textures ({011}<122> and Cube {001}<100>) strongly influences the drawability of Al
alloys. A higher fraction of deformation texture forms ears 45° to the rolling direction
(RD), while a higher fraction of annealing texture forms ears 0°/90° to the RD. Besides
manufacturing processes, processing methods such as severe plastic deformation (SPD),
dissimilar rolling, and accumulative roll bonding can also be employed to tailor the
microstructure and texture [18–20]. The subsequent sections discuss the differences in deformed and recrystallized
microstructures, along with particle simulated nucleation (PSN) and texture evolution,
for various FCC alloys. In this review, the authors provide insight into the generation
of various microstructures and textures that evolve during the deformation and annealing
treatment of FCC materials (such as Al and Cu alloys, ASSs, and HEAs). TMP, a process
widely used to enhance the properties of alloys, results in the formation of a fine-grained
structure, thus enhancing mechanical and chemical properties. Therefore, the authors
discuss the impact of TMP on deformation and texture evolution in several FCC metals
and alloys.
2. ALUMINUM ALLOYS
The deformation and recrystallization behavior of Al alloys is strongly influenced
by the presence of precipitates or intermetallic particles. Achieving optimal deformability
in Al alloys necessitates controlling the microstructure, especially the size of second-phase
particles. The second-phase particles, which include α-Al(Mn,Fe)Si, β-Al(Mn, Fe)Si,
θ-phase, Al6Mn, Mg2Si, Al2Cu, S-phase (Al2CuMg), η-phase (Mg(Al,Cu,Zn)2), and Al7Cu2Fe, influence the mechanical and chemical properties of Al alloys [21–25]. Nondeformable second-phase particles (size>1 mm) create strain incompatibility
at the matrix-particle interface, leading to significant problems during formability
operations. Processing routes define the shape and morphology of the second-phase
particles. DC cast Al alloys contain intermetallics in the form of a Chinese script
[26], whereas TRC exhibits center line segregation (CLS) [27,28]. For TRC AA8011, it has been noted that homogenization at ~580 ºC for 12 h can reduce
the size of intermetallics in CLS, though complete elimination has not been reported
[29].
Fig 2 illustrates the differences in the microstructure and second-phase particles in DC
cast and TRC Al alloys. Equiaxed grain morphology was observed for DC cast AA3000
alloys (Figs 2(a)). For instance, the average grain size (GSavg) after DC cast was 153±47 μm for AA3003. Intermetallics (Chinese script) formed along
the grain boundaries (GBs), as shown in Figs 2(b). During grain formation, Fe/Si/Mn were ejected due to their low solubility in Al
(Fe solubility in Al is 0.05 wt.% at 650 ºC, Si is 1.65 wt.% at 577 ºC, and Mn is
1.8 wt.% at eutectic temperature [21,23]), forming a Chinese script of Al-Fe-Mn-Si along the GBs [30]. Because of the continuous morphology of the intermetallics, it is difficult to
measure their exact size. On the other hand, coarse elongated grains were formed in
TRC due to the simultaneous occurrence of solidification and deformation (Figs 2(c)) [18,27]. Solidification occurred from either side of the roller towards the center of the
cast strip. Therefore, Fe and Si, ejected at the center and due to their low solubility
in Al, formed CLS directly in the center along the production direction, as shown
in Fig 2(d) [27]. CLS, also continuous in nature, consists of numerous branches of intermetallics
along the transverse direction. Consequently, changes in the casting methods alter
the types of microstructure and the formation of intermetallics in the Al alloys.
Of course, a change in texture is expected for both the DC and TRC routes since the
formation of grains varied under both conditions.
Fig 3(a) shows the φ2=0°, 45°, and 65° constant sections of orientation distribution functions (ODFs) for
the FCC metals/alloys. Deformation texture in Al alloys is composed of β-fiber, which
is a continuous tube of texture components that connects the Copper component to the
Brass component through the S, while recrystallization texture consists of Cube and
Goss components. However, the intensity of the textures depends on various factors,
such as % deformation, strain rate, temperature/time of annealing [27], and intermetallic particle size [24]. Fig 3(b) shows the φ2=0°, 45°, and 65° constant sections of ODF for DC-cast AA3003, with the volume fraction
of texture components given in Table 1. A weak q- and g-fiber, along with random orientations were observed in the case
of the DC-cast alloy (Fig 3(b)) [26], whereas strong Copper and weak Goss were observed in TRC AA8011 due to deformation
during solidification, as shown in Fig 3(c) [27]. The q-fiber consists of Rotated Cube and Cube components, while the g-fiber consists
of <111>//ND.
2.1. Deformed microstructure and texture evolution
Deformation in Al alloys increases the stored energy (SE) due to the increment of
dislocation density [31]. Differences in deformation routes lead to variations in the evolved microstructure
(see Figs 4(a-f)). After hot rolling (95% HR) at 500 ºC and 72% RTR of DC-cast AA3003 [26], the grains transform into a lamellar/banded form (elongated in the rolling direction
(RD)), as depicted in Figs 4(a,c) [26]. Additionally, there are unindexed regions (black regions) due to high dislocations
and intermetallic particles. The Chinese script also gets fragmented and scattered
throughout the Almatrix along the RD after the RTR, as demonstrated in Figs 4(b,d). For instance, the intermetallic size varies from ~ 0.2 μm to ~ 10 μm after HR and
RTR [26].
Further, it has been reported that the presence of coarse intermetallics of Al-Mn-Si
(after casting) causes grain fragmentation, exhibiting intragranular misorientation
after severe deformation. The coarse intermetallics also get fragmented along the
RD during 98.84% RTR and refined to the size of ~ 0.2 to ~ 4 μm [32]. Conversely, after TRC, there is no need for additional processes such as HR, unlike
with DC-cast alloy, since TRC produces a strip with a thickness of approximately 7-10
mm that does not require heavy deformation. However, RTR is used to deform the material
up to the required thickness, for example, 2 mm in the case of utensil applications.
During RTR, the elongated grains and center line segregation of TRC AA8011 were further
deformed, leading to the development of a banded structure of grains and fragmented
intermetallic particles, as shown in Figs 4(e, f) [27]. Additionally, refinement of the elongated grains and a decrease in the CLS (the
branches of the CLS come together) were observed [27].
In addition to RTR, equal channel angular pressing (ECAP) has also been used to tailor
the microstructure. Pokova et al. [18] tailored the microstructure of TRC AA3003 via ECAP and, after four instances of
ECAP, elongated grains transformed into ultra-fine grains with a GSavg of 0.5 μm. The transmission electron microscope (TEM) images also confirmed that
the ECAP process evolved subgrains of 500 nm [18]. Therefore, it can be concluded that the HR/RTR/ECAP deformation route leads to
the development of a banded/lamellar type of grain morphology and the fragmentation
of coarse intermetallic particles along the RD. Additionally, plastic deformation
increases the stored energy (SE) by enhancing dislocation density, which will aid
in the recrystallization of new grains during subsequent annealing. With changes in
microstructure, the crystallographic texture also undergoes significant alterations,
dependent on the chosen deformation route. Figs 5(a-c) display the constant sections (φ2=0°, 45°, and 65°) of ODFs for two processing routes: DC cast (Figs 5(a,b)) [26] and TRC (Fig 5(c)) [27], across different Al alloys (AA3003, AA5XXX, AA8011) and post-processing methods
(HR, RTR). The volume fraction of texture components is presented in Table 1. After deformation (HR/RTR), the rolling texture components, such as Brass, Copper,
Goss, and S, are strengthened, while the Cube component decreases, as depicted in
Figs 5(a,b) for the DC cast [26,32]. Similar strengthening of deformation texture components (Brass, Copper, S) was
observed after 72% RTR of TRC AA8011, as displayed in Fig 5(c). This means the rolling texture components were strengthened with a weakening of
the Cube component during HR/RTR deformation routes, which was independent of the
casting route. However, the magnitude of each component varied [26,27,29,32] based on the strain rate, strain, and the HR temperature, as reported in Table 1.
The balance between rolling texture and recrystallization texture (Cube) is crucial
for optimal deformability/drawability. Thus, the Cube component must be present after
deformation to evolve again through strain-induced boundary migration during subsequent
annealing [33]. For instance, up to ~1.7 % and ~1.3 % of Cube was retained after RTR of the DC-cast
(Fig 5(b)) and TRC (Fig 5(c)) samples, respectively, as provided in Table 1.
2.2. Heat-treated microstructure and texture evolution
After deformation, the material becomes significantly work-hardened due to an increase
in dislocations, making it difficult to form any product from the deformed sheets.
Therefore, an annealing heat treatment is required to develop strain-free grains.
Fig 6 illustrates the microstructure and texture development during the annealing treatment
of DC and TRC processed Al alloys. In the DC cast, after annealing at 450 ºC for 16
h, the deformed banded structure was transformed into equiaxed grains of ~24 μm (Fig 6(a)) [26]. The driving force for recrystallization was the stored energy of deformation. After
annealing, the amounts of Fe, Si, and Mn in the intermetallics also increased, forming
stable α-Al12(Fe, Mn)3Si with a size range of ~6-8 μm [26], as shown in Fig 6(b).
Higher deformation (ECAP) can result in a higher SE in the material, producing more
fine grains during annealing. Sidor et al. [32] deformed AA5XXX up to ~98.84% (thickness reduction). In TRC, refinement of grains
up to 14 μm was observed after the annealing treatment at 550ºC for 12 s (Fig 6(c)). Due to the higher SE, recrystallization occurred at several sites rather than growth
of the grains (i.e., nucleation competed with the growth of the grains) during annealing,
resulting in the formation of finer grains [34].
In comparison to the DC cast microstructure, the annealing treatment of TRC AA8011
(72% RTR) at 375 ºC for 50 min displayed the evolution of strain-free, fine grains
with a size of approximately 22±10 μm as shown in Fig 6(d). However, previous studies reported that CLS remained unaffected at 375 ºC and that
the intermetallic size was in the range of approximately 0.63 - 3.63 μm (away from
the CLS) [27]. One study by Kumar et al. [29] focused on the dissolution of the second phase particles during the homogenization
of TRC AA8011. They concluded that the particle size of center line segregation was
reduced to approximately 2 μm at 580 ºC, but there was no complete removal of centerline
segregation. ECAP is an effective method to refine grain sizes. 4-ECAP processed AA3003,
annealed at 400 ºC for 8 hrs, showed a fine grain size of approximately 0.5±0.2 μm,
which is considerably lower compared to other processed Al alloys [18]. Also, TEM confirmed that the ECAP resulted in the evolution of subgrains of 500
nm, which recovered during annealing at 400 ºC.
Moreover, TEM confirmed the presence of Al(Mn,Fe)Si intermetallics in TRC AA3003 that
evolved during preannealing at 450 ºC [18]. After reviewing the microstructural development during annealing, it can be concluded
that the evolution of GS during annealing depends on the amount of deformation and
annealing temperature. A higher amount of deformation enhances the SE and results
in the evolution of fine grains. Moreover, high-temperature annealing reduces the
recrystallization time. Recrystallization texture, mainly consisting of Goss and Cube
orientations, evolves at the expense of deformation texture.
Figs 7(a-c) show constant sections of ODFs at φ2=0°, 45° and 65° after annealing at different temperatures and processing conditions.
These conditions include 450 °C for 16 h (after DC-AA3003+HR+RTR) [26], 375 °C for 240 min (after TRC+72% RTR) [27] and 580 ºC for 8 h (after TRC+50% RTR) [29]. Despite the different processing conditions, in most cases, the Cube orientation
was enhanced during annealing, as shown in Figs 7(a-c) and Table 1. However, the Cube orientation was found to deviate along the ND by 15-20º due to
the instability of the non-octahedral slip system ({110}<011>) during RTR. As a result,
the NDRotated Cube orientation evolved adjacent to S orientation from the deformed
Cube bands [23,35]. Goss-oriented grains also evolved during recrystallization from the Brass orientation
[36]. The % volume fraction of Cube orientation was found to differ for different processing
conditions, which can be attributed to different strain accumulation associated with
the processing conditions, GS, and intermetallic particle size. The control of texture
is essential to achieve good deformability/drawability.
3. COPPER ALLOYS
Electrolytic tough pitch (ETP) and oxygen-free high conductivity are two common grades
of pure Cu [37,38]. ETP Cu is electrolytically refined to a purity of 99.96%-99.99%, with the intentional
addition of oxygen (100-600 ppm), and is known for its high electrical conductivity
(96%-101% IACS) [39,40]. The intentional addition of oxygen forms metal oxides (Cu2O), which react with other impurities to clean the Cu matrix and enhance its electrical
conductivity. ETP Cu possesses excellent electrical conductivity and good formability
[41] but has low tensile strength [42]. Conversely, the addition of solute elements in Cu enhances mechanical strength
and electrical conductivity. Copper ferroalloys (Cu-Fe) have garnered significant
interest due to their excellent magnetic and electromagnetic shielding properties
[43]. Metastable d'-(Cu,Ni)2Si precipitates were formed in the early stages of aging in the Cu-Ni-Si alloy, dictating
the overall properties of the alloys [44]. Rolling deformation is a common technique for creating Cu alloy sheets, which serve
as raw material for the further processing of designated products. Both RTR and cryogenic-temperature
rolling (CTR) have attracted interest as emerging SPD methods for obtaining high work-hardening
and ultra-fine grain microstructures [45,46]. RTR and CTR of Cu and Cu alloys can result in various microstructures, such as
grain elongation, formation of DTs, SBs formation, and kink microstructure [47–49]. RTR-processed Cu-Zn alloy samples (75% RR) have shown an increase in dislocation
density, mechanical twinning, and micro-strain [50]. In Cu alloys, the deformation sequence involves the formation of equiaxed cells
of dislocations, microbands, clustering of microbands, and either SBs or strain localization
formation [51].
3.1. Deformed microstructure and texture evolution
The deformation microstructures reported for Cu alloys are quite different from those
in Al alloys, as, in this case, deformation only starts after forming the cast billet
structure. In this section, the authors attempt to discuss the temperature effect
on microstructure evolution. Deformation at cryogenic temperature (CT) restricts dynamic
recovery and grain growth during deformation, maintaining higher dislocation density
compared to RT-deformed (RTR) samples [52]. Differences were observed when comparing the evolved microstructure in RTR and
CTR pure Cu [53].
Variations in GS, kernel-average misorientation and workhardening were observed for
RTR and CTR processed samples [54,55]. As shown in Fig 8, kernel-average misorientation values were higher for CTR (3.33°) than RTR (2.91°)
up to a 40% reduction ratio (RR), indicating a higher work-hardening rate for the
CTR samples. Further deformation leads to a decrease in kernel-average misorientation
values, which is known as work softening [53]. Work softening infers self-annealing in the deformed Cu samples when exposed to
RT atmosphere [56]. Variations in work-hardening rate create differences in mechanical properties,
and CTR-processed samples exhibit higher mechanical strength than RTR-processed samples.
Significant differences were observed in the micro-hardness and electrical conductivity
values in RTR and CTRprocessed Cu-Mg alloys [57]. A higher rate of increase in micro-hardness values was observed for CTR samples
compared to RTR samples. A 90% CTR sample showed hardness values of ~240 Hv, whereas
a 90% RTR sample showed ~180 Hv.
For Cu alloys, CTR is more effective for improving mechanical strength than RTR. Based
on the microstructural results obtained in the literature [57], several strengthening mechanisms have been suggested to be involved in Cu alloys
i.e., grain boundary strengthening (σGB), twin boundary strengthening (σTB) and dislocation strengthening (σd). RTRprocessed Cu-Mg samples showed higher electrical conductivity than CTR samples.
The electrical conductivity (irrespective of deformation temperature) decreased with
an increase in RRs because of increased GBs, twin boundaries, and dislocation density
[57]. Compared with samples processed at RTR, CTR samples exhibit a larger number of
grain/twin boundaries and high dislocation density, which were responsible for low
electrical conductivity [57]. Solute addition in Cu alloys was able to increase mechanical strength upon RTR;
however, electrical conductivity decreased due to the lattice distortion caused by
those solute atoms [58].
ECAP is another well-known deformation route to apply SPD on Cu. Cobos et al. [42] reported the evolution of a large fraction of low angle grain boundaries (70%, sub-grain
structure with GBs of 3-15°) after the first ECAP pass for pure Cu. Ultrafine grain
microstructures with GSavg of ~0.46 μm and 0.49 μm were obtained after 8 and 16 passes, respectively [42]. The SFE also affects the deformation mechanism, as low SFE Cu alloys (RTR-processed)
show a deformation sequence involving the formation of stacking faults, DTs, and SBs
[48]. One of the features of ETP Cu is the presence of Cu2O particles. Fig 9 shows the presence of Cu2O particles in both the as-received and RTR-processed ETP Cu. From Fig 9(d,e), it can be observed that even after an 80% RR, the Cu2O particles do not fragment, indicating their hard nature [53]. Most Cu2O particles had a size greater than 1 μm, which could assist in PSN during heat treatment
[53]. Voids were observed in the vicinity of these particles due to the deformation incompatibility
between the Cu matrix and strong Cu2O particles [59].
Texture evolution in pure Cu and Cu alloys can be characterized based on the SFE,
where medium to high SFE materials (pure Cu, Al-alloy) exhibit a Copper-type texture
and low SFE materials (Cu alloys, ASSs, HEAs) show a Brass-type texture [60,61]. In the large-strain rolling of initially ultrafine-grain Cu (GSavg = 0.32 μm) obtained by eight passes in ECAP, a Brass-type texture, rather than a
Copper-type texture, was observed. The formation of DTs was the main cause of the
evolution of the Brass-type texture or transition from Copper-type to Brass-type texture
[62,63]. The volume fractions of the major texture components were: 11.2% Copper, 39% S,
and 34% Brass, which is a typical proportion for a Brass-type texture [63]. Visco-plastic self-consistent (VPSC) simulation results revealed that the development
of a Brass-type texture in pure Cu at very large strains likely resulted from the
activation of {11-1}<112> slip, in addition to the usual {1-11}<110> slip. The <112>
slip was only significant in the ultrafine-grain regime, and only a small part (about
10%) of the partial dislocations formed DTs [63]. The evolution of Copper, Brass, and S components has been reported in 80% of RTR
Cu-3Ag-0.5Zr alloys [64]. Similarly, RTR Cu-Cr-Zr alloys showed Goss/Brass ({110}<115>), Brass, and S components
in the 60% RR sample, the intensity of which increased with an 80% RR [65].
3.2. Heat-treated microstructure and texture evolution
High SE and microstructural defects in deformed Cu alloy samples accelerate the recrystallization
kinetics even at lower annealing temperatures [66,67]. Fan et al. [67] reported that an annealing temperature of 300 °C for spin-deformed Cu-Sn alloy was
regarded as the critical point of static recrystallization (SRX). At this temperature,
static recovery occurred, whereas SRX was observed in the annealing treatment from
400-600 °C. SE is gradually reduced at the GBs, leading to the formation of small
grains near GBs through the process of SRX [67]. The microstructure evolution during the recrystallization and grain growth of CTR
Cu was investigated at a temperature of 450 C [68]. Primary recrystallization caused the evolution of very fine grains ranging from
1.59-21.65 μm with many annealing twins. However, long time (100 hours) annealing
at 450 °C caused grain growth with bimodal grain size distribution, in which smaller
grains ranged from 3.35-13.20 μm, with larger grains from 14.96-35.79 μm [68]. Along with high temperature annealing phenomena, self-annealing phenomena have
also been reported for RTR and CTR of pure Cu/Cu alloys [53,69,70]. Self-annealing refers to the occurrence of recrystallization at RT [52]. Lapeire et al. [45] observed the evolution of recrystallized grains in CTR-processed ETP Cu when the
samples were removed from -17 ºC for metallographic preparation. A similar type of
observation was reported by Konkova et al. [71,72]. They also observed a time-dependent softening phenomena in CTR-processed pure Cu.
High defect density, such as vacancies and other lattice defects, can arise during
deformation at CT, which decreases the thermal stability of pure metals and shows
GB migration at RT [73]. Recrystallization phenomena are divided into two types: (1) discontinuous SRX (DSRX)
and (2) continuous SRX (CSRX). Differences in the recrystallization phenomena can
be understood based on the area of nucleation sites and recrystallization texture.
Self-annealing was observed in 80% RTR Cu alloy in which GSavg increased from 57 nm to 280 nm after 1 month of RT exposure [74]. During self-annealing, nanosized/submicron grains were recrystallized in the unique
matrix of a single Brass-oriented deformed grain (after 14 months) [74].
Self-annealing phenomena can be better understood using microstructural explanations.
Differences between the deformed and partially recrystallized (self-annealed) grains
in a CTR80 sample occurred based on the grain orientation spread (GOS) criteria [45,71]. The GOS criterion (2°) was used to differentiate the deformed and self-annealed
grains. Grains with GOS>2° are deformed grains and grains with GOS>2° refer to self-annealed
grains. The ODF of the deformed microstructure showed a plane strain texture (Fig 10(a)). Plane strain texture refers to the presence of Copper, Brass, and S components
which evolved during RTR deformation in Cu alloys. In the magnified region of the
IPF, only deformed grains are shown, while the black partitioned regions indicate
the locations of self-annealed (SRX) grains, which are marked with an arrow and discussed
in Fig 10(b). A high orientation gradient of 12° was also observed inside the deformed grains
(along line 1 in the magnified region in Fig 10(a)), which could cause the further formation of new grains inside the deformed grains
during the process of selfannealing. Interesting results were observed in the case
of the self-annealed microstructure. Fig 10(b) shows the IPF map of SRX grains, predominantly elongated with an GSavg of less than 1 μm. A GOS of > 2º was used to differentiate the SRX grains from the
deformed microstructure. The ODF map in Fig 10(b) illustrates the evolution of plane-strain texture in the SRX grains. The magnified
region of the IPF map shows that some of the SRX grains were elongated, while some
were equiaxed in shape. Gerber et al. [75] studied the microstructure and texture evolution in heat-treated pure Cu after 70%
and 90% reduction in thickness (RTR). Annealing twins were observed, which affected
the growth of the Cube nuclei after 70% rolling and heat treatment (300 °C). In contrast,
the growth of Cube grains was so fast after 90% strain that nothing else was detected
in the electron backscatter diffraction (EBSD) measurement of the recrystallized sample
(150 °C) [75].
Recrystallization texture components, such as Cube and Goss, have been reported for
Cu and Cu-alloys [45,76]. As previously discussed, variations in the SE or RRs can also cause differences
in texture evolution. In deformed conditions (70% and 90% RR), a β-fiber texture with
strong S and Brass components and a weaker Copper component was observed. However,
after 70% rolling and heat treatment (300 °C), both strong Cube and retained rolling
texture components were observed in the ODF map, whereas only the Cube component was
visible in the ODF map after 90% rolling and heat treatment (150 °C) [75]. The effect of SE was also observed; the low RR (70% RR) sample was completely recrystallized
in the temperature range of 200-300 °C, whereas the severely deformed (90% RR) sample
showed an occurrence of recrystallization at 120-150 °C [75]. Of course, severe deformation induces a large SE inside the deformed material,
which enhances the recrystallization kinetics.
Chen et al. [77] reported the evolution of deformation and recrystallization texture in 99% RTR Cu-Ni
alloys. The RTR Cu-Ni alloy showed the evolution of a very strong S component, with
medium Copper, Brass, and weak Goss components. However, the SRX phenomenon at 700-800
°C led to the evolution of a strong Cube texture, whose intensity increased with an
increase in the annealing temperature. Copper, Brass and Goss orientations disappeared
during the annealing treatment [77].
The evolution of PSN microstructure has also been reported for pure Cu and Cu-alloys
[53,78]. Fig 11 shows SEM micrographs of the Cu2O precipitate particles (Fig 11(b)) and their orientation map (Fig 11(c)). Self-annealed grains were observed not only at the deformed GBs but also around
the Cu2O particles [79]. This exceptional behavior in ETP Cu is known as the PSN [53]. As discussed in the Al alloys section, PSN has been reported only for particles
larger than 1 μm [53,78]. A nanostructured Cu matrix was achieved due to the occurrence of discontinuous
dynamic recrystallization (DDRX) and PSN mechanisms after 96% RR asymmetric CTR [78]. The authors would like to inform readers that Figs 10 and 11 are the results of experimental work (ETP Cu) performed for the present review, with
detailed discussions about the sample preparation and results given elsewhere [53]. Fig 11(a) show that a large number of Cu2O particles observed in the CTR80 sample were very rigid, and the shapes of these
particles were not deformed even after 80% RR. During the deformation of ETP Cu containing
second-phase Cu2O particles, dislocations bow around the Cu2O particles. If the strength of the particle is less than the force exerted by dislocations,
the particle will deform; otherwise, the dislocation acquires a semicircular configuration
[80]. The dislocations then encircle the Cu2O particles, leaving an Orowan loop. The generation of extra dislocations, known as
geometrically necessary dislocations (GNDs), in the form of the Orowan loop at Cu2O particles, enhances the SE around the Cu2O particles [81].
Figs 11(b,c) show an SEM micrograph and its corresponding EBSD scan, respectively. Self-annealed
grains can be seen in Fig 11(d); many grains nucleated around the Cu2O particles (marked with a rectangular box). A magnified region near the Cu2O particle in Fig 11(c) displays the nucleation of small grains around the Cu2O particles, as seen in Fig 11(e). PSN (grains 1-10) was confirmed with the help of the GOS≤2º criteria, as shown in
Fig 11(f). The orientation relationship between the PSN grains nucleated around the Cu2O particles was also analyzed. As seen in the (100)-pole figure, grains 1-5 show different
orientations to each other, which indicates the DSRX mechanism (Fig 11(g)) [82]. Similarly, grains 6-10 show dissimilar orientations, indicating that the nucleations
around the Cu2O particles were always DSRX in nature (Fig 11(h)). Generally, PSN causes recrystallization in a discontinuous manner, as also reported
for the Al alloys [82].
4. AUSTENITIC STAINLESS STEELS (ASSs)
ASSs are some of the most widely used FCC materials due to their excellent properties,
including non-magnetic behavior, excellent corrosion resistance, and good formability
and weldability [83,84]. ASS of both the grade 200 and 300 series are metastable at RT, and during deformation,
transform into strain induced martensite (SIM) [6,85]. The 200 series ASSs have a lower SFE compared to the 300 series ASSs, and therefore,
differences in the deformation behavior can be observed between these two grades of
alloys. During deformation, two types of SIM can be formed in ASS: ε-martensite (HCP)
and α'-martensite (BCC) [86]. It has been reported that ε-martensite usually forms at lower strains, reaches
its maximum volumetric fraction, and then decreases with further increases in the
RR [87,88]. For instance, ε-martensite was formed at low strains and reached its maximum volumetric
fraction at ε = 0.11. Further deformation decreases its fraction, whereas the α'-martensite
fraction increased with an increase in the RRs for 201 ASS [86]. As can be understood, the deformed microstructure of ASSs consists of new phases,
i.e., α' and e martensite along with the deformation induced products. Formations
of SIM during deformation can occur via one of the following routes: (a) γ→ ε, (b)
γ→ α', and (c) γ→ ε→ α', which occurs via the displacement of atomic planes [89,90].
In the case of heat treatment/annealing treatment, grain refinement occurs, and SIM
transforms back to a fine reverted austenitic structure (α'→ γ). SIM exhibits a higher
strength level compared to γ, enhancing the strengthening of the steel. Thus, the
mechanical properties, especially the strength of the steel, are mainly controlled
by the formation of α', contributing to pronounced strain-hardening. As discussed
for the previous two materials (Al and Cu alloys), where an improvement in mechanical
properties was attributed to the TMP, was also the case for ASSs. During the TMP process,
the amount of SIM formed during deformation, and refined γ grains formed after heat-treatment,
controlled the overall properties of the ASSs. Therefore, considerable attention must
be paid to estimating the evolution of SIM and refined γ grains during TMP.
4.1. Deformed microstructure and texture evolution
The formation of ε or α' martensite principally hinges on factors like chemical composition,
GS, SFE, and degree of deformation [91]. It has been documented that a decrease in SFE hinders dislocation slip, which is
a primary deformation mechanism, thus creating a window for the occurrence of mechanical
twinning or martensite formation upon straining. 201 ASSs exhibit a notably low SFE
value (approximately 10 mJ/m2), and it has also been reported that an SFE greater than 20 mJ/m2 inhibits ε formation, while a lower SFE leads to the γ→ ε→ α' transformation [86]. Similarly, during the deformation of 316L ASSs, SBs, SIM (ε and α'), DTs, and deformation
bands were formed [90,92]. The typical microstructure evolution of cold-rolled 316LN ASSs is presented in
Fig 12 via optical micrographs (subsets a-d) and TEM micrographs (subsets e-g) [85]. The initial microstructure consists of equiaxed γ-grains (Fig 12(a)). As can be seen, upon a small RR of 10%, SBs (Fig 12(b)) and DTs (Fig 12(e)) were formed, and dislocations were accumulated in them. With an increase in the
RR to 30%, SIM was formed, and the martensite boundary inhibited dislocation movement,
causing dislocations to cluster around the martensite boundary. Upon a further increase
in RR to 50%, the untransformed γ structure was elongated to the RD (Fig 12(c)), and martensite laths were also formed (Fig 12(f)). At 90% RR, a large block of γ and a fine martensite phase mixed with the γ structure
(white circle in Fig 12(d)), and dislocation cell-type martensite was observed (Fig 12(g)) [85].
Zhang et al. [93] investigated the microstructure evolution of cold-rolled and tensile-deformed Cr-Mn-Ni-N
metastable ASS samples at temperatures of 0 °C, -15 °C, and -30 °C. The formation
of α'-martensite was observed during tensile testing, which was sensitive to temperature
but insensitive to RRs. Lowering the deformation temperature effectively promoted
the formation of ε-martensite, which becomes a transitional phase of the subsequent
transformation to α'-martensite, leading directly to the enhancement in strength [93].
For a more detailed view, Fig 13 displays a schematic diagram of the various microstructural features that evolved
with increasing strains [94]. At low strain, dislocation slip predominates, and twinning is uncommon because
the critical resolved shear stress for slip is lower than that for twinning [94]. This microstructure evolved as a result of dislocation glide/slip and typically
forms in alloys with low SFEs. The dislocation cell-block structure is created in
high SFE materials like Al, where dislocations exhibit a high three-dimensional mobility
and can easily cross-slip. Low strain also reveals domain boundaries and microbands
(Fig 13(a)). Microbands and domain boundaries are geometrically necessary barriers created to
account for changes in orientations between adjacent domains. Compared to low strain,
medium strain results in higher dislocation density, finer dislocation border spacing,
and a larger twin volume fraction (Fig 13(b)). Consequently, these boundaries interact more frequently, and the twin-matrix (T-M)
lamellae become an important component of the microstructure. The twinning planes
of the T-M lamellae rotate closely to the rolling plane under high strain. With locally
concentrated slip, deformation becomes heterogeneous, resulting in SBs that are roughly
30° inclined to the RD (Fig 13(c)). The majority of the plastic strain is carried by these SBs, which also aid in martensite
nucleation. It should be emphasized that the scale, alignment, and local strain of
SBs are markedly different from those of microbands (Fig 13(c)).
As discussed in the previous sections, a Brass-type texture was observed for low SFE
materials. The SFEs of 304L ASS, high Mg twinning-induced plasticity steel, and 316L
ASS were 18 mJ/m2, 40 mJ/m2, and 64 mJ/m2, respectively. Chowdhury et al. [95] reported that increasing the RRs for 316L ASS increased the volume fraction of DTs
and SBs. The material initially showed a Copper-type texture, but with the increase
in RRs, the Copper component diminished, and the Brass and Goss components increased.
The mechanism of texture evolution from Copper component to Brass component in medium
SFE materials has been reported by Hirsch et al. [96]. The following routes were observed during the texture transition: Copper→ Copper
twin ({552}<115>) → Goss→ Brass components. Chowdhury et al. [95] reported that the formation of DTs was correlated with the texture transition from
Copper→ Brass-type texture transition. A change in the rolling direction or strain
path can lead to changes in the microstructure and texture evolution in ASSs. The
effect of strain path on the texture evolution of the γ and α' phases was reported
in [97]. Texture evolution for the γ and α phases during unidirectional rolling (RTR) and
the multistep cross rolling of 316L ASS alloy samples was discussed in [97]. Unidirectional rolling refers to samples being placed in the same direction at
every rolling pass, whereas multistep cross rolling indicates that the samples were
rotated by 90° at every rolling pass. Unidirectional rolling resulted in Brass, Goss,
and γ-fiber texture, while multistep cross rolling mainly formed Brass texture for
deformed austenite after 90% RR. The Copper component was not observed in the unidirectional
and multistep cross-rolled samples. After 90% RR, {112}<110> and Rotated-Cube {001}<110>
were the main texture components observed for the transformed martensite in 90% RR
samples of unidirectional and multistep cross rolled samples [97]. On the other hand, during 60% RTR of 201 ASS, the austenite phase developed Goss,
Brass, and S texture components, which did not change significantly upon further straining,
and α' developed Rotated Cube, α- and γ-fibers [86]. Furthermore, the formation of DTs resulted in the development of the Brass component
[86].
4.2. Heat-treated microstructure and texture evolution
As discussed in Section 4, plastic deformation induces SIM in metastable ASS, and
reverting back to the γ phase (α′→ γ) through heat treatment serves as a viable option
for grain refinement [98]. The α′→ γ reversion transformation upon heat-treatment has been explored by various
authors [99,100], with the starting and ending temperatures of this reversal being dependent on the
composition [99]. The reverted γ, forming through the TMP, inherits high densities of dislocations
from the prior α'. Consequently, during reversion annealing, the dislocations reorganize,
establishing finely micro-structured cells and subgrains in the reverted γ phase.
Fig 14 displays the image quality (IQ) maps of SA and RTR (15%, 30%, and 50% RR) samples
after thermal aging (TA) at 900 °C for 6 h for 202 ASS alloy [101]. The IQ micrographs depicted in Fig 14 indicate that as the degree of RRs increases, the GSavg decreased. The GSavg values were 74.081 μm, 48.095 μm, 16.445 μm, and 5.546 μm for the SA, 15%, 30% and
50% RTR, respectively, after thermal aging. The GSavg notably decreased at a rapid rate with higher degrees of deformation (30% and 50%
RRs) after thermal aging, confirming the conclusion of martensitic reversal into smaller,
refined grains of γ, as discussed in [101]. Fig 14(a) reveals only γ grains with large GS, while Fig 14(b) demonstrates that with an increase in deformation to 15% RR, large grains of retained
γ (those γ grains which do not transform into SIM during deformation) are formed with
a fraction of SIM and small reverted γ grains [94,102,103]. At 15% RR, a significant number of dislocations developed near the GBs, with only
a small portion of the grain surface being covered by SIM. This leads to more GBs
becoming active, resulting in sensitization, which elevates corrosion and reduces
mechanical strength [104]. For 30% RR (Fig 14(c)), more refined and equiaxed grains of reverted γ are formed with a very low volume
fraction of SIM, resulting from the dissolution of α' or, put differently, the reversal
of α' into γ [101]. In Fig 14(d), at 50% RTR, grains fully recovered into reverted and refined γ, reducing the GSavg from 100 m to less than 10 mm. It is noteworthy that an increase in deformation escalates
the internal SE of the ASSs, thereby enhancing recrystallization kinetics and potentially
facilitating a more rapid reversion. Diffusional α'→ γ reversion is understood to
primarily be induced by the nucleation and growth of fine grains at martensite lath
boundaries, as observed in 16Cr-10Ni and 18Cr-9Ni metastable ASS samples [105].
In another study, Sun et al. [106] reported that continuous heating of 80% RTR 304 ASS at 700 °C, using heating rates
of 2 °C/s, 20 °C/s, and 100 °C/s, did not significantly alter the texture components.
For continuous heating to 700 °C at 2 °C/s, the γ phase texture was most pronounced
in the Brass orientation, followed by the Goss and Copper components. When the heating
rates were increased, the most prominent texture shifted to the Goss component, followed
by Brass and Copper components. The authors concluded that the heating rate exerted
no significant influence on the annealed γ phase texture [106]. Variations in texture components also became apparent in relation to the fluctuations
in the GSavg of ASS samples [107]. In one study, an as-received 201 ASS alloy sample underwent tensile deformation
(ε=0.34) and subsequent annealing from 100 °C to 800 °C [108]. In the asreceived state, the γ phase displayed a random texture. Upon tensile deformation
(ε=0.34), the Goss and weak Copper texture components emerged in the γ phase, while
in the SIM (e=0.34), α- and γ-fiber, alongside the strong Rotated Cube texture component,
were formed. Upon annealing at up to 590°C after deformation, the Goss component in
the γ phase weakened, while new components, namely Brass, were formed; concurrently,
the texture of the SIM weakened until it disappeared. Notably, the ODF of the reversed
γ phase exhibited only the Brass and S components. Thus, it can be deduced that the
new γ phase nucleated with different orientations. After the complete reversal of
the γ phase, it exhibited a randomized texture [108]. The evolution of texture during annealing at temperatures ranging from 600 to 1000
°C in a 95% RTR 304L ASS was studied, with major components centered around the Goss
and Copper components, as well as the BR component {236}<385>. Entirely new orientations
after the recrystallization of the γ phase were observed, correlating with deformed
texture components through twin relationships. However, a decrease in texture intensity
was observed concurrent with an increase in annealing temperature [109].
The evolution of texture can be linked to deformation texture through twin relationships.
TMP resulted in refined and reverted γ grains, significantly enhancing mechanical
properties in terms of both strength and ductility. Somani et al. [110] demonstrated that the strength of the RTR 301LN steel (and also 301 steel) was exceptionally
high, reaching levels of 1600 to 1800 MPa at high RR, albeit with very low elongation.
However, upon undergoing the reversion process, the strength diminished, while ductility
sharply increased beyond an annealing temperature of 600 °C. Depending on the annealing
temperature (700 to 900 °C), the yield strength ranged from approximately 600 to 1000
MPa, while the elongation varied from 27 to 52% [110].
5. HIGH ENTROPY ALLOYS (HEAs)
HEAs are characterized by the alloying of more than four elements, typically in equiatomic
or near-equiatomic compositions. Despite their complex chemical compositions, HEAs
crystallize in simple crystal structures (FCC, BCC, and HCP) and form non-ordered
solid solutions. High mixing entropy is beneficial for increasing the stability of
the solid solution against the formation of intermetallic compounds [111,112]. Just as in any other conventional alloy, the mechanical and chemical properties
of HEAs are predominantly determined by the microstructure that evolve during TMP.
The trade-off between strength and ductility in HEAs can be addressed by meticulously
designing TMPs to modify the microstructural features to deliver the desired properties
[113–116]. Consequently, a significant field of research has been dedicated to understanding
microstructure and texture evolution in HEAs during various TMP [117]. A fine-grained microstructure can be achieved in HEAs by applying medium strain
(50-70% RR) and annealing at relatively low homologous temperatures (~500-600 °C)
or through a short-time annealing treatment [118]. This finegrained microstructure suppresses DTs, resulting in higher yield strength
but relatively low elongations [119]. SPD leads to ultrafine-grained HEAs, which, when exposed to temperature ranges
of 500-700 °C, result in the formation of intermetallic compounds [120,121]. The degree of deformation prior to annealing (in the temperature range of 700-1000
°C) impacts the recrystallization kinetics and its fraction. Higher strain imposed
before annealing provides more nucleation sites and a larger recrystallized fraction
at a given annealing temperature [122]. Moreover, a fully recrystallized microstructure showed a weak texture, which was
attributed to annealing twinning [111,123–126], which does not depend on the initial GS [127]. Additionally, microstructures with varied proportions of special grain boundaries,
such as coincident site lattice boundaries, can be obtained through different TMP
treatments, which is often referred to as grain boundary engineering (GBE) [116,128,129]. The aforementioned descriptions about the microstructural evolution in HEAs during
deformation and heat-treatment are discussed in the subsequent sections.
5.1. Deformed microstructure and texture evolution
Various microstructural features, including dislocation slip ({111}<110>), DTs ({111}<112>),
and micro-SBs, appear as deformation progresses. The deformation mechanisms can be
controlled by managing the GSavg [118] and SFE [130]. Grain refinement suppresses the formation of DTs, as illustrated by the complete
absence of DTs in an ultrafine-grained CoCrFeMnNi alloy (GSavg=503 nm) due to the very high critical resolved shear stress required for twinning
[119]. Shahmir et al. [131] produced HEAs with GS ranging from 0.05 μm to 70 μm and reported a critical GS of
2 μm below which DTs were not formed. Otto et al. [15] discussed the effect of deformation temperature and grain size on the tensile properties
of the CoCrFeMnNi alloys. The equiatomic CoCrFeMnNi microstructure with coarse grains
(GSavg=157 μm) resulted in a lower yield strength and higher elongation compared to the
fine-grained (GSavg=4.4 μm) microstructure at various temperatures ranging from -196°C to 600°C. DTs
were observed only at -196 °C after a strain of 20%, whereas dislocation cell structures
were developed at RT and high temperatures during tensile loading, regardless of the
GS [15]. In a similar type of study, tensile and compression experiments on a single-crystal
CoCrFeMnNi alloy also supported the fact that DTs only appear at -196 °C and not at
RT deformation [132,133]. In contrast, formation of DTs in the same alloy with GSavg=17 μm were reported prior to necking during a tensile test at RT [134]. DTs were found to be responsible for the increased strain-hardening of the coarse-grained
(GSavg=590.2±9.3 μm) non-equiatomic Fe41Mn25Ni24Co8Cr2 alloy, whereas those features were not observed during deformation of the fine-grained
microstructure sample (GSavg=8.1±0.52 μm) [135]. Furthermore, Co-rich HEAs (Co35Cr20Mn15Ni15Fe15) exhibited higher DTs as compared to the Cantor alloy on account of its lower SFE
(~11 mJ/m2), resulting in higher strain hardening capabilities of Co-rich HEAs compared to that
of the Cantor alloy [136]. In contrast, a highly coarse-grained (GS in the range of 500-1000 μm) non-equiatomic
HEA (Co11.3Cr20.4Fe22.6Mn21.8Ni23.9) also demonstrated good hardness (3 GPa) along
with a good strain-hardening rate (~2600 MPa/ε) [137].
Moreover, DTs appeared in the RTR CoCrFeMnNi alloy (initial GSavg=81±39 μm) after a 20% RR (RR20), and the twin density increased with subsequent RRs
of 40% (RR40) and 60% (RR60) [138], as shown by electron channeling contrast imaging micrographs in Fig 15. When comparing the crystallographic texture with the DTs evolution, the Copper component
was more favorable for the formation of DTs than other texture components. Severe
deformation (RR60) can also cause the formation of micro-SBs and the breakdown of
the dislocation cell structure via DTs, a common phenomenon in low-SFE materials at
higher deformation strains [124,138,139].
The texture evolution of a CoCrFeMnNi alloy (with GSavg of 35 μm) during cryogenic tensile loading was characterized by a predominant {111}<112>
texture component, accompanied by a minor Rotated-Cube ({001}<110>) texture component.
Additionally, the {115}<552> texture component evolved due to twinning of the {111}<112>
texture component [140]. As shown in Fig 16, the typical texture transition from a Copper-type texture to a Brass-type texture
during the RTR of HEAs is consistent with that of other FCC materials with low SFE
[111,138,141]. A similar transition in crystallographic texture was discussed for the rolling
deformation of ASSs in previous sections [95]. The Brass-type texture is characterized by strong α-fiber components, such as Brass
({110}<112>) and Goss ({110}<001>) texture components [138]. The texture that evolved after 90% RR through cryo-rolling (CTR) of CoCrFeMnNi
was not significantly different than the RTR texture, as both exhibited a typical
Brass-type texture [124]. Sathiaraj et al. [127] reported that the initial GS had no significant effect on the texture evolved after
90-95% RTR. When subjected to shear loading during high-pressure torsion experiments,
CoCrFeMnNi alloys (with GSavg of 500 μm) exhibited a texture similar to that of other FCC materials, with major
texture components such as {111}<112> and Copper orientations [125].
5.2. Heat-treated microstructure and texture evolution
The annealed microstructure of HEAs often exhibits numerous annealing twins, which
are introduced due to multiple generations of annealing twinning. In HEAs a lowtemperature
annealing treatment causes the occurrence of static recovery, along with the formation
of intermetallics/precipitations. Zheng et al. [142] conducted annealing experiments at various temperature conditions on cold-rolled
CoCrFeMnNi with different reductions and reported the recovery (<600 °C), recrystallization
(600-830 °C) and grain growth (>830 °C) regimes. Another similar research reported
the recovery temperature to be around 650 °C and that grain growth occurred at a temperature
above 800°C through grain boundary migration in the Cantor alloy [143]. Furthermore, Chen et al. [144] reported a two-stage recrystallization process in the Cantor alloy at 700 °C wherein
the first stage (annealing time<5 min) was characterized by a transition from a cold-rolled
microstructure to a mixture of eutectic-like microstructure and a new, non-recrystallized
microstructure with a changed texture. The second stage was marked by complete recrystallization
[144].
The generation of multiple twinning has also been observed in the Cantor alloy after
80% RTR followed by annealing at 700 °C for 1 h [123,145]. A quasi in-situ annealing experiment performed on the Cantor alloy after RTR revealed
that recrystallized grains preferentially nucleated at SBs and grew by subgrain coarsening
mechanisms [146]. Similarly, a reduction in area of 80% through rotary swaging produced SBs, that
acted as nucleation sites during subsequent annealing treatment [127]. The fraction of twin boundaries was found to depend on the final annealed GS during
the grain growth regime (annealing temperature: 800 and 1000 °C for various times)
for 90% RTR CoCrFeMnNi [147]. In another study, the evolution of annealed microstructures and crystallographic
texture were discussed when the Cantor alloy was subjected to 80% RR followed by annealing
treatment for various time periods, such as 700 °C for 6 min (Fig 17 (a,f)), 700 °C for 7.5 min (Fig 17 (b,g)), 700 °C for 10 min (Fig 17 (c,h)), 700 °C for 15 min (Fig 17 (d,i)), and 700 °C for 1 h (Fig 17 (e,j)) [145]. The annealing treatment led to the evolution of the Copper component and devolution
of α-fiber texture for SRX grains. In the partially annealed samples, non-recrystallized/retained
deformed grains belonged to Brass, Brass/Goss, and Goss orientations [145]. The heterogeneous microstructure underwent hierarchical nano-twinning and showed
higher evolution rates of GNDs. Moon et al. [116] demonstrated an interesting processing route in which the DTs fraction was enhanced
by pre-stretching the Cantor alloy at CT (-196 °C) and then exposing it to a low annealing
temperature of 500°C, which only recovered dislocations while DTs were retained.
The effect of TMP on the microstructural and mechanical properties of carbon added
Cantor alloy was studied by Stepanov et al. [115].The microstructure showed improved strain hardening due to the addition of carbon
(0.2 wt.%) to the Cantor alloy [115]. The twin boundary fraction in the annealed microstructure was lower than that of
undoped Cantor alloy [115]. The increase in SFE due to carbon addition was found to be responsible for this
behavior. On the other hand, the addition of Si (0.2 molar ratio) slowed down the
recrystallization process by hindering GBs migration [148]. Consequently, an increase in Si content led to a more heterogeneous microstructure
with a mix of recrystallized and non-recrystallized grains. Annealing treatment after
plastic deformation promotes the formation of intermetallic phases in HEAs. For example,
prolonged heating of a nano-crystalline Cantor alloy at 450°C resulted in the decomposition
of a single FCC phase into NiMn, FeCo, and Cr-rich phases, which increased the hardness
[120]. In another study, a 95% RTR Cantor alloy followed by annealing at various temperatures
showed the formation of the σ-phase in the microstructure, which further increased
hardness compared to the deformed microstructure [121]. Texture evolution in the Cantor alloy during heat-treatment was similar to the
annealing behavior of other low SFE materials. Crystallographic texture evolution
in the 50% RR Cantor alloy followed by annealing treatment at various temperatures
is mentioned in [111]. It was observed that annealing treatment up to 600 °C, resembles the features of
the deformed texture. However, a further increase in the annealing temperature up
to 700 °C, decreased the texture intensity [111]. The multiple generations of annealing twins introduced many new orientations in
the microstructure and caused texture weakening during the recrystallization of HEAs,
which can be seen for 800 °C and 900 °C [123]. The net effect was significant texture weakening rather than complete randomization.
The effect of strain paths, specially unidirectional rolling and multistep cross rolling,
on the annealing texture of CoCrFeMnNi alloy was studied by Reddy et al. [149]. The dependence of annealing texture evolution on CoCrFeMnNi alloy processed through
different strain paths showed that the {236}<385> texture component did not evolve
during the annealing of cross-rolled samples [149]. Differences in crystallographic texture after annealing were attributed to variations
in the annealed GSavg. The evolution of the microstructure (GSavg) during annealing treatment was dependent upon substructure destabilization and misorientation
build-up, which changed with the varying strain path. Additionally, a higher fraction
of nucleation sites (SBs, GBs, DTs) leads to easier nucleation of recrystallization.
In such cases, nucleation will dominate over the growth process, leading to a finer
recrystallized GS. The deformation structure and misorientation build-up were affected
by cross-rolling, which diminished the density of potential nucleation sites and adversely
affected the nucleation of recrystallization [149]. Annealing of 50% and 80% RTR Al0.25CoCrFeNi was marked by the evolution of the Rotated-Cube component [122]. The intensity of the Rotated-Cube decreased as the annealing temperature increased
from 700 to 1000 °C [122]. In the case of the 90% RTR MnFeCoNiCu alloy, the Brass and Goss components were
found to be quite stable when the annealing time was increased from 1 to 16 hrs at
900 °C [150]. This evolved texture was attributed to the sluggish diffusion effect and the absence
of oriented nucleation and growth during recrystallization. Furthermore, various GBE
treatments, often a combination of deformation and thermal treatments, can also be
employed to increase the fraction of CSL boundaries in the microstructure for enhanced
mechanical and corrosion properties [116,128,129].
6. DISCUSSIONS
6.1. Particle stimulated nucleation (PSN)
Work-hardening and softening are the two main phenomena that occur during deformation
and annealing, respectively, for FCC metals/alloys. Work-hardening refers to the increase
in the fraction of dislocation density during RTR deformation, enhancing the strength
of the materials [151], whereas PSN, recovery, and recrystallization processes lead to softening during
annealing treatment. During deformation, the first aspect to change is the shape of
the grains. Equiaxed grains transform into a lamellar/banded structure, enhancing
the GBs area, which is accompanied by an increase in dislocation density. Hence, the
SE of deformation is the summation of newly formed interfaces and the accumulation
of dislocations [34]. The enhancement of dislocation density can be measured using EBSD in terms of GND
density. GNDs are the dislocations required to accommodate plastic deformation and
maintain compatibilities between the GBs. GND can be influenced by both deformation
temperature and deformation rate/strain rate [152]. Zheng et al. [152] studied microstructural evolution in hot-deformed (tensile testing) AA6082 alloy
samples. It was observed that decreasing the deformation temperature and increasing
the strain rate resulted in increased GND density, leading to higher SE after deformation
[152].
This SE is potential for recrystallization during annealing. In Al alloys, intermetallic
particles (> 1 μm) play a vital role in defining the microstructure and crystallographic
texture through PSN during recrystallization [153]. Fig 18 illustrates the PSN mechanism under the influence of coarse intermetallic particles.
After casting, grains are in an equiaxed shape with coarse second phases, as shown
in Fig 18. Further plastic deformation causes the breakdown of intermetallics distributed along
the RD and results in higher SE. The hard particles induce particle deformation zones
around them [152]. These particle deformation zones consist of higher dislocation density (also referred
to as GNDs) and a large misorientation gradient, indicated in Fig 18 by black dotted lines. Moreover, the size of particle deformation zones is directly
influenced by particle size [154]. The size of particle deformation zones is almost equal to particle size. Larger
particle deformation zones will have a higher misorientation gradient (near the particle-matrix
interface) that will lead to more randomly oriented grains (PSN) during annealing
[153]. This accumulated strain around the particles will be the supporting force for recrystallization
during subsequent annealing. Initially, the dislocations arrange themselves as subgrains
during annealing and then nucleate the randomly oriented grains adjacent to the interface
of the particles and matrix, as shown in Fig 18 [153]. The higher the accumulated strains around the particles, the more random grains
will form [154,156]. Similar PSN behavior was also reported for ETP Cu, which occurred due to selfannealing
as well as high-temperature heat treatment around Cu2O particles [79,157]. In the case of 7Mo super-ASS, s precipitates promoted dynamic recrystallization
through the PSN mechanism [158]. Other literature have also reported the occurrence of the PSN mechanism during
the heattreatment of ASSs [159,160].
6.2. Self-annealing and Softening phenomena
The self-annealing phenomenon depends not only on material properties such as purity,
melting point, and composition but also on external factors such as strain, strain
rate, SFE, and deformation temperature [161]. Under high strain and strain rate and low SFE and deformation temperature, self-annealing
can occur at RT [161]. This phenomenon, observed as SRX and grain growth in Cu, Pb-Sn, Zn-Al, and Al-Cu
alloys, occurs shortly after SPD processing [161,162]. In Al alloys, self-annealing can be correlated with the natural aging process,
which involves the nucleation of very fine precipitates. During natural aging, the
size and number density of the Guinier-Preston (GP) zones in Al alloys change with
time. In contrast to softening at RT in Cu alloys, natural aging enhances hardening
in Al alloys [163]. Softening behavior in Cu alloys (medium to low SFE materials) is divided into SRV
and SRX. As it is well known, a large number of dislocations are generated during
plastic deformation, mainly accumulating around the GBs and particle interfaces. The
difference in the mechanisms of SRV and SRX can be observed in terms of dislocations
arrangement. Of course, SRX occurs at a higher temperature than the SRV process. In
the SRV phenomenon, the following processes of dislocation arrangement can be observed:
(a) dislocation tangles, (b) cell formation, (c) annihilation of dislocations within
cells, (d) subgrain formation, and (e) subgrain growth. Similar types of features
regarding dislocation arrangement during the heat treatment of Cu-Cr and Cu-Cr-Mg
alloys are explained in [164]. The microstructure evolution in Cu-Cr alloy samples after aging treatment at 480
°C showed a cellular substructure after 15 min of heat treatment, while the annihilation
of dislocations and the formation of annealing twins were observed after 4 hrs of
heat treatment [164]. When comparing the softening phenomena in Cu-Cr-Mg alloy samples with Cu-Cr, a
difference in the dislocation arrangement was observed. A large dislocation density
could be observed in the short-timeaged sample. Dislocation tangles were observed
in the 15 min heat-treated sample, whereas a cellular substructure was present even
after 4 hrs of heat treatment [164]. The good softening resistance performance of Cu-Cr-Mg alloy was due to the pinning
effect of dislocations by fine precipitates and Mg atoms [164]. The addition of Ca, Sr and Y was reported to enhance softening resistance in Cu-Cr
alloys [165,166].
6.3. Phase transformation
Phase transformations in Al, Cu and HEAs are unheard of during deformation [26,27,45,138,167]. Rather, annealing or aging treatment at respective temperature ranges causes the
formation of other phase precipitates and/or intermetallics in Al, Cu, and HEAs [120,168–173]. In contrast to the aforementioned FCC alloys, ASSs show phase transformation during
deformation and annealing treatment [89,90]. The TMP of ASS can be performed in two stages: (i) rolling deformation and (ii)
subsequent annealing. As discussed in Section 4, in the first stage, the transformation
can occur via one of these routes: (a) γ→ ε, (b) γ→ α', and (c) γ→ ε→ α', which increases
the yield strength and microhardness of the ASS [85,174]. The study showed that the ε-martensite formed at lower RR; however, by increasing
the RR, the ε-martensite was observed to be coexisting with α'-martensite, and α'-martensite
could transform from the surrounding ε-martensite [174]. Also, it was reported that RTR promotes the formation of α' for various grades
of ASSs [89,92]. Furthermore, as discussed above, the transformation cycling (deformation) results
in an increase in dislocation densities; it has been reported that high dislocation
densities significantly decreases γ-phase stability [175], and by decreasing the stability of the γ-phase, the kinetics of the martensitic
transformation could be enhanced [175,176].
Mohammadzehi et al. [176] investigated the effect of RTR and γ-phase stability on the mechanical properties
of 316L ASS. As is known, 301 and 201L ASSs have lower γ-phase stability compared
to 316L ASS. The lower the γ-phase stability, the higher the chances of martensitic
formation, even at low RRs. γ-phase stability depends upon the alloying content in
ASSs; lower alloying contents in 301 and 201L ASSs compared to 304 and 316 ASSs effectively
mask the effects of other variables such as varied GS and small changes in the rolling
temperature. In the second stage, a reversion of the α′→ γ phase takes place, resulting
in improved ductility, grain refinement, and enhanced corrosion properties. Proper
execution of TMP leads to a fine-grained γ structure, with GSavg ranging from nano to submicron, imparting excellent room-temperature strength and
ductility to the ASSs [177,178]. Plastically deformed metastable ASSs can undergo two types of reverse transformation
during TMP: an athermal ‘shear-type’ transformation and an isothermal (thermally activated
'diffusional') reverse transformation [179]. Both transformation types lead to γ-phase formation, but the resulting microstructures
are completely different. The thermally activated isothermal transformation is time-dependent
and leads to equiaxed, finely grained γ-phase with low dislocation density [179]. In contrast, the athermal transformation is temperaturedependent only and leads
to a lath-shaped γ-phase with high dislocation density [180]. The type of reverse transformation experienced by an ASSs depends on its chemical
composition and the heating rate applied during the heat treatment [181,182]. Higher heating rates and temperatures enable athermal transformation, whereas low
γ stability enables isothermal transformation [183].
The reversion mechanism of α′→ γ during thermal aging at 700 °C at different heating
rates in the experimental steel is given in [106]. The study concluded that the diffusional reversion process dominated the transformation
of α′→ γ at low heating rates (2°C/s), resulting in the development of equiaxed, defect-free,
nano/ultrafine-grained austenite grains. While at a rapid heating rate (100°C/s in
this work), the regulated process was a martensitic shear-type reverse transformation
of α′→ γ that produces a lath-type reversed γ structure with a high density dislocation
[106]. Further, Mao et al. [175] reported that the two step rolling and annealing process was beneficial in achieving
ultra-fine grained ASSs which exhibited a great combination of strength and ductility.
The study reported that in a metastable Fe-24Ni-0.3C alloy sample, the first step
rolling and annealing process (50% RTR + 600 °C annealing for 30 s) resulted in a
sample consisting of a partially recrystallized γ-phase with GSavg ~1.5 μm, and upon the second step rolling and annealing process (90% RTR + 600 o°C
annealing for 30 s), the sample consisted of a fully recrystallized γ-phase with GSavg ~0.5 μm [175]. Fine-grained ASS possesses better mechanical properties than coarse-grained ASS
[107].
6.4. Grain boundary engineering
As discussed in Section 5, the concept of GBE is to control the fraction of special
boundaries in the microstructure, such as CSL boundaries in low SFE FCC materials.
Twin boundaries (<111>60°) in FCC materials belong to Σ3 CSL boundaries. The fact
that FCC alloys are compliant with GBE imparts distinct mechanical and corrosion behaviors
to the material due to the low GB energy of these special boundaries [116,128,129,184,185]. Chen et al. [128] reported that a CSL boundary fraction of over 50% was achieved during a one-step
recrystallization of CoCrFeMnNi. It has been reported that the fraction of annealing
twin boundaries decreases with increasing amounts of alloying elements. This is because
the alloying effect and severe lattice distortion can decrease both the average GB
energy and twin boundary energy [128]. However, Thota et al. [129] argued that a proper GBE treatment should disrupt the network of random HAGBs and,
therefore, strain-annealing (low strain followed by annealing) should be preferred
as a GBE treatment over the one-step recrystallization employed by Chen et al. [128].
The effect of annealing temperature and time on CSL boundary fraction for non-deformed
(as-received) and RTR Cantor alloy samples was discussed in [129]. The authors reported an exceptional CSL boundary fraction (>70%) in HEAs obtained
through 5% RR (RTR) followed by 1 h annealing at 950 °C when compared to other conditions,
such as, as-received conditions, 10% RR and 15% RR (RTR) followed by 1 h annealing
at 950 °C. Moon et al. [116] adopted a strategy to enhance the twin boundary fraction in HEAs by deformation
at CT, followed by low-temperature annealing that was just sufficient for the recovery
of the microstructure without affecting the twin boundaries. Moreover, Kaushik et
al. [123] also reported a CSL fraction of more than 50% in Cantor alloy, which was RTR to
80% RR, followed by annealing at 700 °C for 1 hr. The microstructure consisted of
several twin clusters with profuse Σ3 and Σ9 boundaries. Higher RTR deformation and
annealing temperatures increased the twin boundary fraction in the HEAs [186]. For the case of hot deformation of ASSs, it was concluded that the generation of
strain-free grains by dynamic recrystallization and the occurrence of CSL boundaries
(Σ3 and Σ9) led to a random texture [187]. In another study on RTR ASSs [94], DTs occurred on planes with the highest twinning Schmid factors and showed a strong
orientation dependence. In the case of RTR 316L ASS, DTs occurred preferentially in
grains with near Copper orientation rather than the Brass orientation [94].
7. SUMMARY AND RECOMMENDATIONS
In the present review, TMP and its impact on microstructure and texture evolution
have been thoroughly studied for Al and Cu alloys, ASSs, and HEAs. Various features
of the deformed microstructure in FCC materials, such as grain-fragmentation, precipitates/particles
distribution, SBs/DTs formation, and martensitic transformation, are discussed for
the aforementioned alloys. Similarly, texture evolution during deformation and annealing
treatment and the effect of second-phase particles, strain rate and deformation temperature,
annealing twin boundaries (CSL boundaries), and phase transformation on evolved crystallographic
texture have been discussed for various grades of FCC metals/alloys. A concise summary
has been drawn for each alloy case as follows:
1. DC-cast and TRC Al alloys exhibit second-phase particles in the form of Chinese
script and CLS, respectively. RT deformation fragments coarse intermetallic particles
of Chinese script/CLS into smaller particles. Low-temperature (<450 ºC) annealing
evolves a greater fraction of secondphase particles, whereas at high-temperature annealing
(> 500 ºC), the size of intermetallics is reduced. PSN have been reported only for
the coarse intermetallic particles (size > 1 μm) during annealing treatment. However,
detrimental effects of PSN have been observed, as it reduces the likelihood of forming
preferred orientations (especially the Cube {100}<001>) after heat treatment.
2. Grain fragmentation and the formation of SBs/strain localizations and DTs have
been reported during the plastic deformation of Cu alloys. Severe deformation enhances
the SE of Cu alloys, which causes self-annealing or RT recrystallization phenomena.
Softening occurs when the specimen is exposed to the RT atmosphere. The formation
of SRX grains has been observed at the deformed GBs and triple junction boundaries,
which is due to the DSRX mechanism, whereas the formation of grains inside the deformed/parent
grains is due to the CSRX mechanism. PSN has also been reported in ETP Cu, which causes
the nucleation of grains around the hard Cu2O particles via the DSRX mechanism. Pure Cu exhibits a Copper-type texture, whereas
low SFE Cu alloys exhibit the Brass-type texture after RTR deformation.
3. In ASSs, the nucleation characteristics of SIM (phase transformation) during TMP
were observed to occur via (a) γ→ ε, (b) γ→ α', or (c) γ→ ε→ α'. During subsequent
heat-treatment, recovery occurs, through which SIM reverts back to refined γ grains.
However, the process of recrystallization can only occur in the reverted γ regions
when α'→ γ transformation is complete. Upon deformation, the texture developed in
304L is a Copper-type texture, while 316L develops only the Brass-type texture. The
texture transition during severe deformation (from Coppertype to Brass-type) in medium
SFE materials follows the route: Copper component→ Copper twin ({552}<115>) → GossBrass.
4. In single FCC phase HEAs, the evolution of DTs depends on GS, deformation temperature,
and the state of stress. Larger GS, low deformation temperature, and deformation by
rolling greatly promote twinnability during deformation behavior. In contrast, smaller
GS, temperatures higher than RT, and low strain tensile/compression loading impede
deformation twinning. The deformation texture exhibits a transition from the Copper-type
texture to the Brass-type texture with an increase in RTR reduction. The SBs developed
during severe deformation act as preferred nucleation sites for recrystallized grains.
Complete recrystallization results in the evolution of a weak texture due to the introduction
of several new orientations as a result of multiple annealing twinning in HEAs.